Optimization of microstructure and mechanical properties of microwave-sintered V/Ta/Re-doped tungsten heavy alloys

3.1.1 XRD analysis

The XRD analysis results in Figure 2 reveal the phase composition characteristics of the five tungsten alloys (S1–S5) prepared by microwave sintering. These alloys are dominated by W phase, with Ni-Fe-Co phase. From the spectrum, the five curves from bottom to top correspond to S1 to S5, with the body-centered cubic W phase dominating. The characteristic peaks are located at 2θ ≈ 40.3° (110), 58.3° (200), 73.2° (211), and 87.0° (220). The intensity and half-width of these peaks reflect the grain size and strain distribution of W. S1, as the baseline alloy, shows clear W peaks with the highest intensity, along with the face-centered cubic (111) peak (2θ ≈ 43.6°) and (200) peak (2θ ≈ 50.6°) of Ni, the hexagonal close-packed (104) peak (2θ ≈ 36.3°) of Co, and the weak (200) peak (2θ ≈ 65.3°) of Fe, indicating the formation of a stable binder phase of Ni-Fe-Co during microwave sintering at 1,470 °C. This promotes liquid-phase diffusion and fills the gaps between W particles, resulting in higher density. This is consistent with traditional powder metallurgy, but the volumetric heating effect of microwaves accelerates liquid-phase formation and suppresses W grain coarsening. After introducing 1 wt% V in S2, the XRD spectrum only retains the W, Ni, and Co peaks from S1, with no peaks detected for V or related compounds. This suggests that V atoms incorporated into the W matrix via solid solution, forming a W(V) solid solution rather than a distinct phase [23]. The low V content may be below the XRD detection limit, thereby maintaining phase purity and potentially enhancing solid solution strengthening of W. The addition of 2 wt% Re in S3 introduces a (212) peak of ReW, indicating that Re partially forms a σ phase or ReW compound with W, promoting grain boundary pinning and grain refinement, which can improve the alloy’s tensile strength and ductility [3]. Furthermore, the weak ReW peak indicates its formation is constrained; the absence of a pronounced independent Re peak suggests that most of Re atoms have been incorporated into the tungsten lattice through solid solution. The addition of 4 wt% Ta in S4 introduces the body-centered cubic (221) and (110) peaks of Ta, reflecting that Ta’s high melting point causes it to maintain an independent phase at the sintering temperature and not fully dissolve in W. S5, which is co-doped with multiple transition metal elements (1V-2Re-4Ta), only shows the four main peaks of W, the (111) peak of Ni, the (104) peak of Co, and the (110) peak of Ta in the XRD, with no independent peaks for V and Re or the ReW peak. This is not a detection flaw but results from a complex interaction mechanism: under the non-equilibrium conditions of microwave sintering, V and Re may preferentially solvate into the W lattice or the Ni-Fe-Co binder phase, forming disordered solid solutions (W(V,Re) or Ni(V,Re)), and the high Ta content (4 wt%) as an “anchoring” phase stabilizes the overall structure, suppressing the nucleation of compounds like ReW. This phase evolution is related to the complexity of Ta-Re co-doping.

Figure 2:

XRD patterns of different tungsten alloy samples.

3.1.2 Surface morphology analysis using optical microscopy

OM observations of microwave-sintered WHAs reveal the regulatory effects of transition metal element doping on W grain morphology, binder phase distribution, and microdefects. For sample S1, Figure 3(a) shows that the W mostly retains its spherical morphology, with Ni-Fe-Co binder phase uniformly filling the particle gaps as a binder, promoting densification through capillary action. However, a small amount of microporosity is visible at the interfaces, resulting from local liquid phase insufficiency or gas escape during the sintering process. These defects are rare and support the foundation of liquid-phase sintering. In S2 with 1 wt% V added (Figure 3(b)), the W grain size slightly decreases, and the spherical profile becomes less distinct. Due to the interface-active adsorption of V, the wettability between W and binder phase is enhanced, promoting particle rearrangement, and suppresses binder phase coarsening. The micropores further decrease, optimizing the density [24]. In sample S3 (Figure 3(c)), the W grain size significantly reduces, and the distribution becomes polycrystalline. The binder phase still continuously bridges the grain boundaries, but the pinning effect induced by Re, while refining the structure and improving hardness, increases local penetration unevenness due to reduced liquid phase viscosity, leading to a slight increase in micropore density. In Figure 3(d), the morphology of S4 is similar to S1, with a slight reduction in W size and retention of spherical features. However, the high melting point of Ta limits its diffusion, forming a localized Ta-enriched phase, and the number of micropores slightly increases due to heterogeneous aggregation that obstructs liquid-phase flow [25]. In S5, multiple transition metal elements synergistically act on W grains (Figure 3(e)), with a grain size similar to S3 but slightly larger. The distribution is more heterogeneous, with the small-grain region predominantly refined by Re/V, while the large-grain region ensures the stable presence of Ta. The binder phase network size is uniform, and the V/Re/Ta interaction balances solid solution and dispersion. The low-density micropores maintain high hardness and density, but the grain differences amplify the stress gradient. Overall, optical microscope observations provide good evidence of the progressive regulation by doping: V/Re refines the grains and increases strength, Ta introduces a hardening effect but slightly increases defects, and multiple transition metal elements achieve comprehensive optimization in sample S5.

Figure 3:

Light microscope images of different tungsten alloy samples: (a) S1, (b) S2, (c) S3, (d) S4, (e) S5.

3.1.3 SEM surface topography analysis

The SEM images in BSE mode reveal the surface microstructure characteristics of microwave-sintered tungsten alloys. Figures 4–6 distinguish phase regions using atomic number contrast: the light gray region corresponds to the high-Z W phase, while the dark gray region represents the low-Z Ni-Fe-Co binder phase. During sintering at 1,470 °C, the liquid phase fills the gaps between W particles through capillary action, promoting densification and maintaining the spherical morphology of the W powder. This reflects the non-equilibrium effect of microwave volumetric heating, avoiding grain recrystallization and coarsening typical of traditional sintering. Meanwhile, the microporosity observed in Figure 4(a) results from local binder phase insufficiency or gas escape, with the micropore count being very low and not significantly reducing density, though they act as stress concentrators, affecting fracture toughness. Figure 4(a1)–(a4) show the EDS surface scan of Figure 4(a), confirming that the W phase is enriched with W, while the binder phase is enriched with Ni, Fe, and Co, with a uniform distribution and no obvious segregation. The boundary line scan in Figure 4(b) further shows that the interface transition zone is 1–2 μm wide, with Ni-Fe-Co enriching the interface and forming a diffusion layer about 0.5 μm thick. This promotes metallurgical bonding, enhances interface adhesion, and suppresses crack initiation, consistent with the classical model of liquid-phase sintering, where Ni-Fe-Co alloying reduces liquid phase viscosity and accelerates W particle rearrangement [21]. In contrast, after introducing 1 wt% V in S2, Figure 4(c) shows that the binder phase size significantly decreases, which is attributed to the interface-active effect of V. V atoms adsorb at the W/binder phase interface, improving binder phase wettability, while also reducing pore formation. This confirms that V preferentially solutes into the W lattice, forming a W(V) solid solution, and strengthens the solid solution through lattice distortion, improving the hardness of the W phase and enhancing ductility [26]. The binder phase composition in Figure 4(c2) shows approximate composition of W:Co:V ≈ 6.75:13.62:8.92 wt% (as determined by semi-quantitative EDS analysis), indicating that V partially dissolves into the binder phase, possibly forming a V-Co microalloyed phase, which further stabilizes liquid phase viscosity and suppresses phase separation. These semi-quantitative EDS results suggest that V partially dissolves into the binder phase, consistent with the observed refinement of the binder phase network.

Figure 4:

Microstructure of samples S1 and S2. (a) Surface morphology of sample S1; (a1)–(a4) show the surface scan results for (a); (b) elemental distribution maps of the binder phase and its boundaries obtained from line scans; (c) surface morphology of sample S2; (c1) shows the V element distribution for (c); (c2) shows the point scan elemental content map for the “1·” region in (c).

Figure 5:

Microstructure of samples S3 and S4. (a) Surface morphology of sample S3, (b) enlarged image of the boxed area in (a), (b1)–(b5) EDS elemental distribution maps of (b), (c) surface morphology of sample S4, (c1)–(c5) EDS elemental distribution maps of (c), (d) line scan elemental distribution map of the boxed area in (c).

SEM backscattered electron images also visually show the surface microstructure characteristics of microwave-sintered tungsten alloys S3 and S4. These images use grayscale contrast to clearly distinguish the W phase (light gray) and the Ni-Fe-Co binder phase (dark gray). In Figure 5(a), the overall morphology of S3 continues the typical two-phase distribution, with the binder phase network uniformly filling the gaps between W particles, while a few micropores are scattered. These pores mainly result from local gas residue or insufficient liquid phase flow during sintering, but the overall density remains high. The magnified region in Figure 5(b) further reveals that the binder phase size is significantly reduced compared to S1, and the spherical profile of W particles becomes increasingly blurred, transitioning towards a polycrystalline structure. This reflects the interface reconstruction induced by Re addition, which enhances particle contact and promotes the fine distribution of the binder phase. The EDS images in Figure 5(b1)–(b5) show the distribution patterns of Fe, Ni, and Co, all of which are highly localized in the binder phase region, forming a continuous bonding network. The signal distribution of W and Re overlaps significantly, indicating that Re is uniformly dispersed in the W phase. The significant overlap of Re and W signals in EDS mapping indicates predominant solid solution of Re in the W phase, consistent with limited ReW compound formation detected by XRD. This dual behavior-primarily solid solution strengthening combined with minor ReW phase for grain boundary pinning-contributes to the observed grain refinement and enhanced hardness. For sample S4, the surface morphology in Figure 5(c) shows that the binder phase network still exists, but the spherical features of W are further weakened, and the particle edges become more blurred, accompanied by an emerging Ta-enriched phase (appearing as island-shaped light gray blocks). This significant heterophase is due to the high doping level of 4 wt% Ta. The Ta particles did not fully dissolve into the binder phase, resulting in local aggregation. Energy-dispersive spectroscopy analysis in Figure 5(c1)–(c5) directly confirms the enrichment of Ta, with the W phase dominating the overall structure. Fe/Ni/Co still anchor the binder phase, confirming the phase separation tendency of Ta. Line scan analysis in Figure 5(d) of the Ta-enriched region and its neighborhood shows that the concentrations of W and Ta are significantly higher than Fe/Ni/Co, especially in the dark gray core region, where Ta content exceeds that of W. This highlights the extreme localization of Ta, likely due to its high melting point limiting its diffusion, leading to the formation of Ta-rich regions under microwave heating [27]. In conclusion, these image observations align with the XRD results. The Re-W overlap in S3 strengthens phase uniformity, while Ta enrichment in S4 introduces a heterophase but may contribute to dispersive strengthening. Furthermore, the observed near-spherical W grains, reduced microporous structure, and uniform distribution of dopant elements are attributed to the rapid volumetric heating and non-equilibrium effects during microwave sintering. These effects suppress the grain coarsening and segregation phenomena typically observed in conventional sintering processes.

Next, the surface microstructure of S5 is presented in detail, highlighting the phase domain complexity induced by multi-transition metal element doping. In Figure 6(a), the W phase boundary is highly blurred, and the spherical profile is almost indistinguishable. The binder phase network exhibits size heterogeneity, with large and small sizes coexisting. This differentiation reflects local regulation by transition metal elements during the sintering process, leading to uneven liquid phase flow and weakening the geometric retention of W particles. The magnified view in Figure 6(b) focuses on the yellow box area and clearly distinguishes the grayscale variation of the binder phase. The light gray phase intermingles with the dark gray phase, accompanied by areas of transition metal element aggregation. These areas range from 3 to 7 μm in size, indicating the aggregation effect of transition metal elements. The elemental surface distribution images in Figure 6(b1)–(b7) directly show significant overlap of the Ta and Re signals with W. Both elements are evenly dispersed in the W phase region, while V and Ta exhibit noticeable local aggregation. The V signal peaks in specific areas, indicating the tendency for V/Ta agglomeration. Fe/Ni/Co, on the other hand, is localized in the small-sized dark gray phase and only extends into a few large-sized light gray phase regions. This confirms that the binder phase network prioritizes small binder phase regions, forming efficient interface bridging. Line scan analysis in Figure 6(c) reveals elemental variations across the large-sized binder phase profile. The W concentration peak region is accompanied by synchronous peaks from Ta and Re. Quantitative point scanning in Figure 6(d1)–(d4) further refines the regional features. In Figure 6(d1), the light gray transition metal aggregation region is dominated by Ta (81.95 at%), with V as a secondary element (14.64 at%) and trace amounts of Re (0.3 at%). W only makes up 1.11 at%, reflecting the V-Ta synergistic effect. In Figure 6(d2), the Ta concentration in the dark gray aggregation region decreases to 63.24 at%, while V increases to 33.02 at%, with darker colors corresponding to higher V density. In Figure 6(d3), the binder phase region shows Fe/Ni/Co peaks (Fe 8.73 at%, Co 10.37 at%, Ni 19.4 at%), with lower concentrations of transition metals and W, which helps improve binder purity. In Figure 6(d4), the W matrix region is nearly pure W (94.84 at%), with trace amounts of transition metals (Ta 1.03 at%, V 1.46 at%), maintaining the uniformity of the matrix. Overall, these image observations reflect the heterogeneous microstructure of S5. The W-Re/Ta overlap strengthens phase uniformity, while Ta enrichment in S4 introduces a heterophase but may contribute to dispersive strengthening.

Figure 6:

Microstructure of sample S5. (a) Surface morphology of sample S5; (b) enlarged image of the boxed area in (a); (b1)–(b7) EDS elemental distribution maps of (b); (c) line scan elemental distribution map of the green circled area in (b); (d1)–(d4) elemental distribution maps of the “1·”, “2·”, “3·”, and “4·” four-point regions in (b). “4·”.

3.2.1 Sintering density

The relative densities of the five WHA samples prepared by microwave sintering are shown in Figure 7. The baseline sample S1 achieves 98.6 ± 0.61 %, which thanks to the efficient filling of the W particle gaps by the Ni-Fe-Co liquid phase formed in the sintering process and the densification promoted by microwave volumetric heating [17]. The relative density of S2 with 1 wt% V added slightly decreases to 97.32 ± 0.42 %, possibly because the interface activity of V improves wettability but introduces local solid solution distortion, slightly increasing the grain boundary porosity. S3 has the lowest value of 95.7 ± 0.60 %. Although Re-induced grain boundary pinning refines the grains, it reduces the liquid phase flow, leading to uneven penetration and an increase in micropores, which is consistent with the increased porosity observed in fracture analysis. The relative density of S4 increases to 97.92 ± 0.72 %. Although Ta, as a high-melting-point dispersive phase, aggregates locally, it enhances the skeletal support and optimizes overall filling. S5, with multiple transition metal elements co-doped, has a relative density of 96.33 ± 0.50 %, which is between S3 and S4. The combined effects of V/Re/Ta balance solid solution strengthening and binder phase network heterogeneity. The micropore size is small and the quantity is low, but the differences in grain size amplify local defects. The overall trend reflects the trade-offs involved in doping with transition metal elements.

Figure 7:

Relative density corresponding to the prepared WHA samples.

3.2.2 Hardness

The Vickers hardness test of microwave-sintered tungsten alloy samples (Figure 8) shows that the hardness of the baseline S1 is 380.92 ± 8.45 HV, mainly due to the inherent hardness of the W phase and the metallurgical bonding of the Ni-Fe-Co binder phase. The hardness of S2 with 1 wt% V added increases to 407.22 ± 6.93 HV. V solutes into the W matrix, forming W(V) lattice distortion, achieving solid solution strengthening and improving uniformity. The hardness of S3 with 2 wt% Re added reaches the highest at 443.27 ± 13.33 HV. Re induces grain boundary pinning, refines the grains, and forms ReW compounds. The dual mechanism significantly enhances deformation resistance, and despite having the lowest density, the strengthening effect dominates. The hardness of S4 (4 wt% Ta) is 434.1 ± 18.91 HV. Ta, as a high-melting-point dispersive phase, contributes to second phase strengthening. However, local aggregation leads to a large standard deviation, and the hardness is slightly lower than that of S3. The hardness of S5 with multiple transition metal elements is 440.25 ± 18.0 HV. The synergistic solid solution and dispersion of V/Re/Ta balance each other, approaching the peak value of S3. However, grain heterogeneity and micropores introduce fluctuations. The overall trend reflects that doping with transition metal elements optimizes hardness, with single Re being the most optimal.

Figure 8:

Hardness values for five tungsten alloys.

3.2.3 Tensile test

The tensile properties of five WHA samples (S1–S5) were analyzed, the values reported represent the average values from three parallel tests (n = 3) for each composition, and the representative tensile stress-strain curves of all samples are shown in Figure 9. S1 has an average tensile strength of 788.25 MPa and an elongation of 8.1 %. As the baseline sample, it shows a good balance between strength and ductility. The Ni-Fe-Co liquid phase effectively fills the gaps between W particles, and microwave sintering suppresses grain coarsening, which is the main reason for its performance. S2 has a tensile strength of 688.2 MPa and an elongation of only 2.51 %. Although V doping provides solid solution strengthening through lattice distortion, the observed significant reduction in tensile strength and elongation compared to S1 indicates that excessive lattice stress and potential local stress concentrations induced by V outweigh its beneficial effects in this composition, leading to embrittlement. S3 has the highest tensile strength of 874.88 MPa, but its elongation is extremely low (0.25 %). Re doping substantially refines grains and forms minor ReW phases, resulting in the highest tensile strength but extremely low elongation, indicative of brittle fracture. The enhanced interface ductility suggested by dimple formation in fracture analysis is insufficient to compensate for the overall brittleness caused by reduced liquid phase viscosity and increased micropores. S4 has a tensile strength of 688.08 MPa and an elongation of 0.78 %. Ta’s high melting point forms dispersive phases to improve hardness. However, local aggregation and phase separation prevent a significant increase in strength and limit ductility [28]. S5 has a tensile strength of 816.35 MPa and an elongation of 8.92 %. The synergistic effect of multiple elements optimizes performance. V solid solution strengthening, Re grain refinement, and Ta dispersive strengthening comprehensively improve strength. The heterogeneity of the binder phase network is balanced, with better control over grain size differences and micropores, and ductility is significantly better than other doped samples [29]. The overall trend shows that Re significantly enhances strength but sacrifices ductility. The effects of V and Ta doping alone are limited. Sample S5, with multi-element doping, achieves a synergistic optimization of strength and ductility, making it more suitable for high-performance applications.

Figure 9:

Tensile strength-elongation curves for five tungsten alloys.

3.2.4 Fracture analysis

Fracture analysis of the five tungsten alloy composites (S1–S5) reveals the different effects of transition metal element (V, Re, Ta) addition and the microwave sintering process on the microstructure. For sample S1, the SEM image in Figure 10(a) shows that the W powder largely maintains its spherical morphology, indicating that the microwave non-equilibrium heating effect suppressed grain recrystallization. The magnified region in Figure 10(b) shows two types of fracture morphology: the rough fracture surface is similar to transgranular fracture, while the smooth surface suggests a mixed ductile and brittle fracture mode. The Ni-Fe-Co liquid phase, acting as a binder, fills the gaps between W particles, promoting densification. The small amount of microporosity at the W/binder phase interface (indicated in (b)) may result from incomplete binder phase penetration or gas residue, potentially becoming stress concentrators. The EDS mapping images in Figure 10(b1–b4) confirm that the distributions of W, Fe, Co, and Ni are uniform. W dominates the spherical phase, and Ni-Fe-Co concentrates in the binder phase, supporting effective metallurgical bonding.

Figure 10:

Tensile fracture of samples S1 and S2. (a) Fracture surface morphology of sample S1; (b) enlarged image of the boxed area in (a); (b1)–(b4) EDS elemental distribution maps of (b); (c) fracture surface morphology of sample S2; (d) enlarged image of the boxed area in (c); (d1) V element distribution in (d).

For the sample S2 with 1 wt% V added, the fracture morphology in Figure 10(c) shows that the spherical profile of W powder is weaker compared to S1, indicating that V has induced grain boundary modification. The magnified image in Figure 10(d) shows that the fracture surface is mainly rough, with the smooth area almost disappearing, suggesting a transition to transgranular fracture, which may be related to solid solution strengthening of V in the W matrix. Microporosity is still scarce, consistent with S1, but it may affect fracture toughness. The EDS image in Figure 10(d1) shows that V is uniformly distributed in the W matrix, with only a small amount of local aggregation. This suggests that V atoms preferentially solute into the W lattice to form W(V) solid solution, which is consistent with the XRD result where no independent V phase was detected. This may be due to the low content and the fast sintering dynamics suppressing compound formation. The reduction in smooth fracture surfaces in S2 indicates enhanced interface bonding and grain refinement, which may improve hardness and tensile strength. The consistent micropore features in both samples suggest that the binder phase content or sintering parameters need to be optimized to reduce defects.

For sample S3, the SEM image in Figure 11(a) shows significant refinement of W grains, with only a few larger grains remaining. This is due to the grain boundary pinning effect induced by the addition of Re. Re atoms either solute or form ReW compounds, inhibiting W particle coarsening and promoting the uniformity of grain rearrangement during liquid-phase sintering. The magnified view in Figure 11(b) further reveals that the fracture surface is almost entirely smooth, with many dimples, indicating a transition to intergranular ductile fracture as the dominant mode. The presence of dimples reflects that Re has enhanced the ductility of the W/binder phase interface. The increased diffusion layer thickness of the binder phase at the grain boundaries can promote plastic deformation rather than brittle fracture. However, the number of pores between W grains slightly increases, possibly because the interface adsorption of Re reduces the liquid phase wettability, leading to insufficient local penetration or gas retention. While these micropores do not significantly affect the overall density, they may serve as crack initiation points, reducing fracture toughness. The EDS mapping in Figure 11(b1) confirms that the Re element highly overlaps with the W grains, supporting that Re preferentially solutes into the W matrix, forming a W(Re) solid solution. This solid solution strengthening through lattice distortion improves the alloy’s tensile strength and high-temperature stability, which is almost consistent with the overlapping Re-W signals observed in the surface SEM.

Figure 11:

Tensile fracture of samples S3 and S4. (a) Fracture surface morphology of sample S3, (b) enlarged image of the boxed area in (a), (b1) Re element distribution map of (b), (c) fracture surface morphology of sample S4, (d) enlarged image of the boxed area in (c), (d1) Ta element distribution in (d).

For sample S4, the fracture morphology in Figure 11(c) shows mixed fracture characteristics, with both smooth and rough fracture surfaces present, and the latter being more dominant. This indicates that the fracture behavior is between intergranular and transgranular modes. The high melting point of Ta limits its complete solid solution at 1,470 °C sintering, causing Ta to partially exist as an independent dispersive phase. This induces local stress unevenness and promotes the formation of rough transgranular fracture surfaces to release energy. The magnified image in Figure 11(d) reinforces this observation, with blurred edges of the rough fracture surface, accompanied by regions of Ta enrichment. These heterophases, while contributing to dispersive strengthening and improving hardness, may also amplify interface weaknesses, increasing crack propagation paths. The presence of smooth fracture surfaces is attributed to the dominant filling of the liquid phase in the non-Ta regions, maintaining some ductility. The EDS image in Figure 11(d1) shows that Ta distribution is uneven, with areas of high concentration. This confirms Ta’s tendency for phase separation, which is consistent with the Ta-W separation observed in the surface SEM line scan. This localization may stem from Ta’s low diffusion coefficient, leading to the formation of Ta-rich cores under microwave volumetric heating, potentially reducing overall uniformity. Overall, the mixed fracture surface of S4 reflects the trade-off of Ta doping: it enhances high-temperature performance but sacrifices ductility. Compared to the uniform refinement in S3, S4 needs further optimization of Ta content to reduce the agglomeration of transition metal phases.

The fracture morphology of sample S5 with multi-transition metal element co-doping is shown in the Figure 12. Its microscopic fracture characteristics reflect the regulation of the complex interactions between V, Re, and Ta under microwave sintering conditions, forming a balance of heterogeneous phases and strengthening mechanisms. Figure 12(a) shows a significant heterogeneity in W grain size, with some larger grains maintaining spherical profiles, while smaller grains tend to become polycrystalline. This can be attributed to the competing effects of various transition metals: Re pinning refines small grains, the high melting point of Ta causes insufficient solid solution in larger grains locally, and the interface activity of V promotes liquid-phase rearrangement, overall suppressing coarsening [30]. The magnified view in Figure 12(b) reveals that the fracture surface is predominantly smooth, with a partial distribution of rough fracture surfaces, favoring an intergranular ductile fracture mode. The dimples are more abundant, likely resulting from binder phase plastic deformation and Re/V dislocation pile-ups, enhancing ductility. The rough surface is associated with Ta-enriched heterophase, promoting transgranular dissipation. Micropores are still present, but their low quantity and small size remain unchanged. These micropores are primarily caused by uneven liquid-phase penetration and gas escape residue during the sintering process. Although they do not significantly reduce density, they can easily act as stress concentrators, triggering crack initiation and propagation, potentially weakening the alloy’s fracture toughness and fatigue life. The EDS spectra in Figure 12(b1–b7) confirm that W dominates the grains, and V, Re, and Ta almost coincide with W, forming a W(V,Re,Ta) solid solution. The lattice distortion strengthens hardness and stability, while Fe, Co, and Ni form the binder phase network, improving interface bonding.

Figure 12:

Tensile fracture of sample S5. (a) Fracture morphology of sample S5; (b) enlarged image of the boxed area in (a); (b1)–(b7) show the EDS elemental distribution maps of (b).